Rapidly solidified intermetallic-second phase composites

ABSTRACT

This invention relates to a rapidly solidified product comprising a second phase in both a stable particulate form and a metastable flake form dispersed in an intermetallic matrix.

This invention was made with Government support under contract NoN00014-86-C-2277 awarded by the U.S. Naval Research Lab. The Governmenthas certain rights in this invention.

CROSS REFERENCE TO RELATED APPLICATION

This application, is a Division of U.S. patent application Ser. No249,626, filed Sept. 26, 1988, now U.S. Pat. No. 4,915,905 which is acontinuation-in-part of U.S. patent application Ser. No. 873,889, nowU.S. Pat. No. 4,836,982, and 873,890, now U.S. Pat. No. 4,774,052, bothfiled June 13, 1986, which are in turn continuations-in-part of Ser. No.662,928, filed Oct. 19, 1984, now abandoned.

BACKGROUND OF THE INVENTION

The present invention relates generally to a process for the rapidsolidification of certain composite materials and to a composite producthaving an intermetallic containing matrix including an in-situprecipitated second phase, such as another intermetallic phase or aceramic material, wherein the second phase comprises a boride, carbide,oxide, nitride, silicide, sulfide, etc., or intermetallic of one or moremetals.

1. Field of the Invention

For the past several years, extensive research has been devoted to thedevelopment of metal-ceramic composites, such as aluminum reinforcedwith carbon, boron, silicon carbide, silica, or alumina fibers,whiskers, or particles. Metal-ceramic composites with good hightemperature yield strengths and creep resistance have been fabricated bythe dispersion of very fine (less than 0.1 micron) oxide or carbideparticles throughout the metal or alloy matrix. However, this technologyhas not extensively been used to produce composites having intermetallicmatrices.

Intermetallics such as titanium aluminides are receiving increasedattention for application as high performance structural materials. Inparticular, these ordered intermetallic compounds offer improved hightemperature properties, including enhanced strength-to-weight ratios andoxidation resistance relative to conventional high temperature titaniumalloys. However, general exploitation of these alloys has been limitedby the lack of significant room temperature ductility and toughness, aswell as the technical challenges associated with processing and/ormachining the material into a final, usable form.

Compositing of titanium aluminides with particulate or fiberreinforcements creates the potential for additional improvements inalloy performance. For example, incorporation of a dispersed phase canresult in direct strengthening of the matrix via dispersion orsecond-phase mechanisms, as well as stabilizing a fine matrix grainsize. The latter can lead to additional improvements in processability(via enlargement of a stable processing "window"), as well as potentialimprovements in strength, ductility, and toughness.

Use of rapid solidification techniques creates the potential foradditional alloying strategies, for example, the incorporation of rareearth alloying additions to produce a homogeneous, nanometer-scale rareearth oxide dispersion. Production of near-net-shape compositeintermetallic components via powder metallurgy (P/M) techniques canminimize fabricability and machining problems inherent to intermetallicalloys.

2. Description of the Prior Art

Prior art techniques for the production of metal-ceramic composites maybe broadly categorized as powder metallurgical approaches, molten metaltechniques, and internal oxidation processes. The powder metallurgicaltype production of dispersion-strengthened composites would ideally beaccomplished by mechanically mixing metal powders of approximately 5micron diameter or less with the oxide or carbide powder (preferably0.01 micron to 0.1 micron). High speed blending techniques orconventional procedures such as ball milling may be used to mix thepowder. Standard powder metallurgy techniques are then employed to formthe final composite. Conventionally, however, the ceramic component islarge, i.e., greater than 1 micron, due to a lack of availability, andhigh cost, of very small particle size materials since their productionis energy intensive, time consuming, and costly in capital equipment.Furthermore, production of very small particles inevitably leads tocontamination of the particles with oxides, nitrides, and materials fromvarious sources. Further, in many cases where the particulate materialsare available in the desired size, they are extremely hazardous due totheir pyrophoric nature.

Alternatively, it is known that proprietary processes exist for thedirect addition of appropriately coated ceramics to molten metals.Further, molten metal infiltration of a continuous ceramic skeleton hasbeen used to produce composites. In most cases, elaborate particlecoating techniques have been developed to protect the ceramic particlesfrom the molten metal during admixture or molten metal infiltration, andto improve bonding between the metal and ceramic. Techniques such asthese have resulted in the formation of silicon carbide-aluminumcomposites, frequently referred to as SiC/Al, or SiC aluminum. Thisapproach is only suitable for large particulate ceramics (e.g., greaterthan 1 micron) and whiskers, because of the high pressures involved forinfiltration. The ceramic material, such as silicon carbide, is pressedto form a compact, and liquid metal is forced into the packed bed tofill the intersticies. Such a technique is illustrated in U.S. Pat. No.4,444,603, of Yamatsuta et al, issued Apr. 24, 1984. Because of thenecessity for coating techniques and molten metal handling equipmentcapable of generating extremely high pressures, molten metalinfiltration has not been a practical process for making metal-ceramiccomposites.

The presence of oxygen in ball-milled powders used in prior art powdermetallurgy techniques, or in molten metal infiltration, can result inoxide formation at the interface between the ceramic and the metal. Thepresence of such oxides will inhibit interfacial binding between theceramic phase and the matrix, thus adversely effecting ductility of thecomposite. Such weakened interfacial contact can also result in reducedstrength, loss of elongation, and facilitated crack propagation. Inaddition, the matrix may be adversely effected, as in the case oftitanium, which is embrittled by interstitial oxygen.

Internal oxidation of a metal containing a more reactive component hasalso been used to produce dispersion strengthened metals, such asinternally oxidized aluminum in copper. For example, when a copper alloycontaining about 3 percent aluminum is placed in oxidizing atmosphere,oxygen may diffuse through the copper matrix to react with the aluminum,precipitating alumina. This technique, although limited to relativelyfew systems since the two metals utilized must have a wide difference inchemical reactivity, has offered a feasible method for dispersionhardening. However, the highest possible level of dispersoids formed inthe resultant dispersion strengthened metal is generally insufficient toimpart significant changes in properties such as modulus, hardness, andthe like. In addition, oxides are typically not wetted by the metalmatrix, so that interfacial bonding is not optimum.

Because of the above-noted difficulties with conventional processes, thepreparation of metal-ceramic composites with submicron ceramicdispersoids for commercial applications has been extremely expensive.

Another class of materials which has seen considerable interest anddevelopment is intermetallic materials, especially intermetallics ofaluminum such as the aluminides of titanium, zirconium, iron, cobalt,and nickel.

The need for the advanced properties obtainable with intermetallicmaterials is typified by their potential application to structurescapable of withstanding high temperatures, such as turbine engines. Indesigning and operating turbine engines today and for the foreseeablefuture, there are two primary problems which demand solutions from thefield of materials science. The first of these is the need to operatecertain portions of the engine at higher gas and metal temperatures toimprove operating efficiency and save fuel. The second problem is theneed for lighter materials to decrease engine weight and engineoperating stresses due to heavy rotating components, and to increase theoperating life of disks, shafts, and bearing support structures. Theselatter structures require materials which are less dense than the nickelbase superalloys they are intended to replace, but which possess roughlythe same mechanical properties and oxidation resistance as thosematerials in current usage.

The intermetallics are typically highly ordered compounds, in the sensethat they possess regularly repeating (e.g., A B A B A B) atomsequencing. Intermetallic compounds are particularly suited to theseneeds because of two properties which derive from the fact that theypossess ordered structures. Modulus retention at elevated temperature inthese materials is particularly high because of strong A-B bonding. Inaddition, a number of high temperature properties which depend ondiffusive mechanisms, such as creep, are improved because of thegenerally high activation energy required for self-diffusion in orderedalloys.

The formation of long range order in alloy systems also frequentlyproduces a significant positive effect on mechanical properties,including elastic constants, strength, strain-hardening rates, andresistance to cyclic creep deformation. Finally, in the case ofaluminides, the resistance to surface oxidation is particularly goodbecause these materials contain a large reservoir of aluminum that ispreferentially oxidized.

However, during metallurgical processing, one problem encountered isthat these materials tend to form coarse grains, which degrade certainmechanical properties, the most important of which is ductility. Also,in many intermetallics the strong A-B bonding results in low temperaturebrittleness, although the exact mechanism of the ductile-brittletransition seems to be different for the different intermetalliccompounds. It is thus necessary to address the problem of minimal lowtemperature ductility without destroying the inherent high temperaturestrength and stiffness. In the prior art it has generally beenconsidered that these latter high temperature properties may only beretained by preserving the ordered structure, hence sacrificing lowtemperature ductility.

Since the early 1970's, the pace of work on ordered alloys andintermetallic compounds has slackened, as a result of lack of progressin improving either ductility or creep resistance of these otherwisevery intriguing alloys.

Interest in utilizing ordered alloys for structural applications wasreawakened in the United States when researchers discovered thatductility and strength improvements could be achieved in TiAl and Ti₃ Albased alloys using a combination of powder metallurgy and alloyingtechniques. Later work on the titanium aluminides utilized ingotmetallurgy. The development of rapid solidification methods led torenewed interest in the iron and nickel aluminides. The replacement ofcobalt in Co₃ V by nickel, and then iron, led to a series offace-centered cubic Ll₂ -type superlattices with greater ductility atambient temperatures. Also, it has been reported in Japan thatpolycrystalline Ni₃ Al can be made more ductile by adding smallquantities of boron. Later, this work was confirmed and the criticalcomposition range over which boron was beneficial was identified. (SeeU.S. Pat. No. 4,478,791 of Huang et al, assigned to General Electric.)These discoveries, together with the national search for replacementsfor strategic metals, such as cobalt and chromium, and the need todevelop energy-efficient systems, have in the past few years stimulatedmuch additional works, largely in the area of improving low temperatureductility and increasing high temperature strength.

An example of work related to intermetallic matrix composites is taughtin U.S. Pat. No. 4,774,052, of which this is a Continuation-In-Part.This patent application teaches a method for forming composite materialsof descretely dispersed particulate second phase materials inintermetallic matrices, particularly in aluminide matrices. Thedispersed material may constitute a second phase such as a ceramic, oran intermetallic compound other than the matrix metal.

Similarly, extensive research and development has been conducted in thearea of rapid solidification (RS) processing. Rapid solidificationprocessing effects highly desired forms of alloys. Homogeneous materialat or above melt temperatures is subjected to a rapid quench ortemperature drop to "freeze" the material to desired microstructure. Therate at which the melt is quenched is in the range of approximately 10⁴° C. per second to 10⁸ ° C. per second. See, for example, U.S. Pat. No.4,402,745, hereby incorporated by reference.

Current technological interest in materials produced by RS processing,especially when used to produce fine powders followed by consolidationinto bulk parts, may be traced, in part, to problems associated with thechemical segregation that occurs in complex, highly alloyed materialsduring conventional ingot casting and processing. During processing viaslower cooling rates used for conventional casting processes, solutepartitioning, that is, macro-and micro-segregation of different alloyphases present in these alloys, and the formation of undesirable,massive particle boundary eutectic phases, can occur. Metal powdersproduced directly from the melt by conventional powder productiontechniques, that is, drop tower, inert gas or water atomization of themelt, are usually cooled at rates three to four orders of magnitudelower than those that can be obtained by RS processing. The latterremoves macro-segregation altogether and significantly reduces spacingover which micro-segregation occurs, if it occurs at all.

Design of alloys made by conventional slow cooling processes isprimarily dictated by the corresponding equilibrium phase diagrams.Alloys prepared by such processes are in, or at least near, equilibrium.The advent of rapid quenching from a melt has enabled divergence fromequilibrium and has added new alloys with unique structures andproperties for commercial use.

Rapid quenching, or rapid solidification, techniques are known formanufacture of metal powder for powder metallurgical (PM) purposes byfinely "atomizing" molten metal. Here, RS occurs not by contact but"in-flight". This technique permits little time for particle growth. Thesmall drops produced solidify to form small granules, each one of whichessentially constitutes an "ingot" of the molten metal. These smallgranules can be charged into a container that is evacuated and sealed.Afterwards, the small granules are compacted and concurrently orsubsequently heated. This compaction and heating joins together thegranules into a solid metal compact of the molten metal composition.This method is valuable for producing homogeneous materials from meltalloys which, if conventionally processed, would result in large-scaleheterogeneities and segregation. Additionally, RS can produce materialscontaining fine metastable dispersoids and second phases.

Prior art techniques for "atomizing" molten metal have includedimpingement, melt spinning, and nozzle atomization.

In impingement techniques, atomization of molten metal into small dropsis usually brought about in inert gas, such as argon or nitrogen. Thegas impinges as high speed jets upon a pouring stream of molten metal.Water and steam have also been used. However, water and steam areunsuitable in certain instances because they cause severe oxidation ofgranules.

It is also known to atomize a pouring stream by impingement onto arotating disk to make small drops or "ingots" which then solidify bycontact with the surrounding atmosphere, cooling-water or oil bath, or acoolant shower. As mentioned above, in this approach the solidificationdoes not occur by contact with the disk. That contact forms the drops orspheres which can have nearly monosize distribution.

British Patent Specification No. 519,624, hereby incorporated byreference, relates to powdered or granular metallic products constitutedof solidified metallic particles derived from molten metal. It alsodescribes a method of producing the product. These solidified metallicparticles have spontaneously crystallized from a metastable undercooledstate at a predetermined temperature below but close to the freezingpoint of the metal. The particles have substantially uniform size andcomposition

To produce such particles, molten metal is discharged from a suitablereceptacle in one or more streams onto a metal surface of such naturethat sufficient heat is abstracted from the molten metal to lower itstemperature to that of an undercooled state, that is, to a temperaturewhich is slightly below the freezing point of the particular metal butwithout causing solidification or crystallization. This surface uponwhich the molten metal impinges can be a belt or a disk rapidly movingeither linearly or rotatively, respectively. The molten metal isimmediately converted into a stream of film-like proportions on thesurface and the extent of the belt or disk surface is such that themolten metal makes contacts therewith for a period just sufficient toundercool it as above defined. Then the molten metal is caused to leavethe supporting surface and to continue its travel in the same directionand at substantially the same speed for a sufficient distance to causesolidification. Because the undercooled stream of film-like proportionshas little or no integrity, it immediately breaks up into a myriad offine, small liquid particles which solidify to form a powdered metal.

These operations may be carried out in a vacuum or suitable atmosphere,and the myriad of fine, small liquid particles may be passed through acoolant to hasten solidification of the particles or to reduce thedistance needed for solidification. During solidification, surfacetension causes the particles to assume a substantially spherical shape.

One known rapid solidification technique involving a centrifugalatomizing process is taught in U.S. Pat. Nos. 4,025,249 and 4,343,750,hereby incorporated by reference. It uses forced convective cooling ofmolten droplets to achieve cooling rates on the order of 10⁵ -10⁶ °C./sec. This rapid solidification state is designated RSR. Such a RStechnique, in conjunction with powder metallurgy techniques forconsolidation of the rapidly solidified powders, has produced materialswith metastable phases, very fine grain structures, highroom-temperature strength and good high temperature properties up to thepoint of instability of the metastable phases.

An approach to further enhance certain material properties is to blendthe RS powder with ceramic powders prior to consolidation. This leads toimprovement in some mechanical properties, for example, modulus,hardness, etc. Silicon carbide aluminum (SiC/Al ), such as commerciallyavailable SiC/7090, produced by an RS/PM approach is an example of sucha material. The difficulty with this approach is that it suffers fromproperty and processing disadvantages inherent to a PM process. Thesedifficulties include a relatively coarse reinforcement (greater than 1micron) and/or weak metal/ceramic interfaces due to surfacecontaminants.

One alternative to conventional RS/PM techniques for developingdispersion strengthened alloys is to form the ceramic dispersoid phaseduring RS processing. U.S. Pat. No. 4,540,546, hereby incorporated byreference, describes a "Melt Mix Reaction" (MMR) process involvingchemically reacting two starting alloys in a mixing nozzle in which amelt mix reaction takes place between the chemically reactablecomponents of starting alloys to form submicron particles of theresulting compound in the final alloy. The mixing and chemical reactionis performed at a temperature which is at or above the highest liquidustemperature of the starting alloys but which is also substantially belowthe liquidus temperature of the final alloy, and as close to the solidustemperature of the final alloy as possible. Whiledispersion-strengthened alloys can be produced by this technique, thereappear to be a number of inherent difficulties. First, processing istechnically complex. Second, efficient mixing is important if finedispersions are to be consistently produced. Lastly, very high degreesof superheat would be required to completely dissolve the RS alloyingelements in order to produce high dispersoid loadings, which necessarilyaccentuate particle growth, for example, in one containing 10-20 %dispersoid.

In U.S. Pat. No. 4,240,824, hereby incorporated by reference, Moskowitzet al describe a process for producing a boron-containing nickel orcobalt spray-and-fuse self-fluxing alloy powder containing an internallyprecipitated chromium boride or nickel boride. In this patent, thestarting materials are alloys containing precursors of the hardprecipitate, and the melt is precooled to a temperature about 50° F.higher than the viscous temperature prior to atomization. The particlesare formed in the secondary atomization step, and are preferably largerthan 10-15 microns in average particle size. No teaching is found forprecipitating the particulate material prior to the atomization steps,or of precipitates having an average size less than 1 micron.

Narasimhan, in U.S. Pat. No. 4,268,564, hereby incorporated byreference, teaches the preparation of sheets or strips of amorphousmetal containing embedded particulate matter, of 1 to 100 micronparticle size, by forcing a glass-forming alloy containing particulatematter, formed in-situ, onto a rapidly moving chill surface. Thistechnique was considered surprising because it had previously beenbelieved that incorporation of particulate matter, especially ofwettable particulate matter, into a molten glass-forming alloy wouldpreclude quenching into an amorphous body due to nucleation ofcrystallization. Further, inclusion of particulate material in the metalmelt in a melt spin process has led to rapid plugging of the orifice.The reference does not teach preparation of a rapidly solidified powderhaving an evenly dispersed particulate material therein. In fact, thereference specifically teaches that the particulate material isconcentrated at the surface of the strip material produced.

These prior art techniques produce conventional powdered metal products.

U.S. Pat. No. 4,836,982, of which this application is aContinuation-In-Part, relates to an invention which overcomes thedisadvantages of the prior art noted above, including current rapidsolidification technology, and provides for the rapid solidification ofcomposite materials comprising metal and metal alloy matrices.

The cited invention may also result in improvement from incorporation ofa stable dispersoid into the composite which extends the hightemperature working range of the composite, in contrast to conventionalRS composites that typically contain metastable phases. Moreover,incorporation of dispersoids prior to RS may provide surfaces forprecipitation, and consequently, a more efficient precipitation ofmetastable rapid solidification phases. In some cases, for example,titanium-based alloys, the addition of rare earth elements, like ceriumor erbium, to the dispersoid-containing melt may result in improvedscavenging of interstitials such as oxygen, leading to the formation ofan additional oxide dispersoid and effective deoxidation of the matrixalloy.

It has now been found that one may form an intermetallic-second phasecomposite in accordance with the teachings of U.S. Pat. No. 4,774,052,and subject such material to rapid solidification in accord with theteachings of U.S. Pat. No. 4,836,982, to achieve a rapidly solidifiedintermetallic matrix of very fine grain size having a second phasedispersed therein.

SUMMARY OF THE INVENTION

This invention relates to a method for the production of rapidlysolidified intermetallic-second phase composite materials. The methodinvolves the formation of a composite material comprising anintermetallic containing matrix having in-situ precipitated second phaseparticles dispersed therein. This composite is subsequently subjected torapid solidification processing to produce a composite powder.

The present invention relates to a rapidly solidified product containinga preformed dispersion of in-situ precipitated second phase particles inan intermetallic matrix, produced by reacting second phase formingconstituents in the presence of an intermetallic in which theconstituents are more soluble than the second phase.

The invention further relates to a rapidly solidified compositecontaining a preformed dispersion of second phase, having a particlesize of from about 0.01 to about 5 microns, in a matrix of anintermetallic in which the constituents of said second phase are solubleand the second phase is substantially insoluble.

Thus, it is a particular purpose of this invention to provide rapidsolidification products containing in-situ precipitated second phases inan intermetallic matrix composite. This approach provides an alternativeto conventional ingot metallurgy or powder metallurgy techniques as ameans of preparing composites for rapid solidification atomization. Theimportance of this distinction resides in the fact that most compositesprepared by conventional techniques are unsuitable for atomization. Theadvantages of the present method and product of this invention willbecome more readily understood by consideration of the followingdescription and examples.

BRIEF DESCRIPTION OF THE FIGURES

FIG. 1 is a photomicrograph showing a cross section of a rapidlysolidified titanium aluminide-titanium diboride composite powderproduced in accordance with the present invention.

FIG. 2 is a photomicrograph showing a cross section of a conventionalunreinforced rapidly solidified titanium aluminide powder possessing adendritic microstructure.

FIG. 3 is a photomicrograph showing the microstructure of a rapidlysolidified titanium aluminide-titanium diboride composite powderproduced in accordance with the present invention wherein a bimodaldistribution comprising equiaxed TiB₂ particulates and fine flake-likeTiB₂ particulates is present within the titanium aluminide matrix.

FIG. 4 is a photomicrograph showing the microstructure of a rapidlysolidified, heat treated titanium aluminide-titanium diboride compositepowder produced in accordance with the present invention wherein abimodal distribution comprising equiaxed TiB₂ particulates and finespherodized TiB₂ particulates is present within the titanium aluminidematrix.

DESCRIPTION OF PREFERRED EMBODIMENTS

The present invention relates to a rapid solidification processutilizing preformed intermetallic-second phase materials, produced by anin-situ precipitation process. Overall, it is the purpose of the presentdisclosure to describe an approach to producing intermetallic rapidsolidification powders having dispersoids therein by a relatively simplemethod. Materials thus produced possess improved mechanical propertiescompared to those produced by known techniques, while maintaining knownadvantages. The method of producing rapidly solidified products of thepresent invention differs from prior art in that at least one dispersedphase pre-exists in the intermetallic product prior to the rapidsolidification treatment. The process thereby circumvents the need foran excessive degree of superheat required in order to dissolve alloyingelements and permits inclusion of higher amounts of dispersoids.Further, the invention avoids the need for controlling excessivereaction exothermicity released during formation of dispersoids in theprior art, and allows greater latitude in selection of dispersoids andreactant concentration.

An in-situ precipitation technique is described in detail in U.S. Pat.No. 4,710,348, hereby incorporated by reference. The preparation ofsuitable intermetallic-second phase composites is described in detail inU.S. Pat. No. 4,774,052, and application Ser. No. 190,561; 190,547 and156,682, hereby incorporated by reference. The specific rapidsolidification technique utilized in the present invention is describedin detail in U.S. Pat. No. 4,836,982, incorporated by reference, filedJune 13, 1986.

Broadly stated, the present invention constitutes the rapidsolidification of novel intermetallic-second phase composite materials.The materials which are subjected to rapid solidification are preparedby a process whereby elements forming a second phase precipitate reactin a solvent matrix material containing an intermetallic, or at leastone precursor thereof, to form a finely-divided dispersion of the secondphase material in the solvent matrix material. While ceramic materialsconstitute the preferred second phase dispersoids, it is also possibleto precipitate intermetallic dispersoids as well. In such instances, thedispersoid and matrix must be of different intermetallic composition.While the discussion herein shall focus upon ceramic materials as thesecond phase, it must be borne in mind that intermetallic second phasesare also to be included in the scope of the present invention.

Exemplary of suitable second phase precipitates are the borides,carbides, oxides, nitrides, silicides, oxynitrides, sulfides, andoxysulfides. Suitable elements include all of the elements which arereactive to form ceramic precipitates, including, but not limited to,transition elements of the third to sixth groups of the Periodic Table.Particularly useful ceramic phase forming constituents include aluminum,titanium, silicon, boron, carbon, oxygen, nitrogen, sulfur, molybdenum,tungsten, niobium, vanadium, zirconium, chromium, hafnium, cobalt,nickel, iron, magnesium, tantalum, manganese, zinc, lithium, beryllium,thorium, and the rare earth elements including scandium, yttrium,lanthanum and the lanthanide series elements such as cerium and erbium.Reactive compounds of such elements, such as B₂ O₃, B₄ C, and BN mayalso be used.

It is especially to be noted that plural dispersoids, and/or complexcompounds such as titanium zirconium boride, may advantageously beprecipitated in-situ in the intermetallic matrix. Additionally,particles having whisker-like morphology may be precipitated, such astitanium niobium and boride.

Composites of relatively low oxygen content may be produced byincorporation of small amounts (e.g., up to about 5 percent, dependentupon oxygen content of the matrix) of strong oxide formers, such asyttrium or any of the rare-earth metals, e.g., cerium and erbium, whichwill scavenge oxygen from the solvent matrix material. The exothermictemperature rise of the reaction mass, in conjunction with the increasedsurface area of the dispersoid formed, may effectively improvescavenging capability. The oxide particles thus formed enhance matrixductility by removal of interstitial oxygen, and may also serve todispersion strengthen the matrix and reduce grain size therein, in turnfurther enhancing matrix ductility. Further, it is to be noted that inmany intermetallic-ceramic composites prepared by the method of thepresent invention, intermetallic precursors will also react with ceramicconstituents to form additional ceramic dispersoids.

As the matrix or solvent metal, one may use an intermetallic, orprecursors thereof, capable of dissolving or at least sparinglydissolving the constituents forming the second phase, and having alesser capability for dissolving or otherwise reacting with the formedsecond phase precipitate. Thus, at the temperatures experienced duringthe process, the matrix component may act as a solvent for at least oneof the second phase-forming constituents, but not for the desired secondphase precipitate. It is important that the second phase-formingreaction goes to substantial completion by consuming the reactants.While the potential choice of second phase dispersoids and matrixmaterials is large, this choice is limited by adherence to the criteriahereinabove recited.

The solvent metal precursors for the intermetallic matrix may beselected from the group consisting of aluminum, nickel, copper,titanium, cobalt, iron, platinum, gold, silver, niobium, tantalum,boron, zinc, molybdenum, yttrium, hafnium, tin, tungsten, lithium,magnesium, beryllium, thorium, silicon, chromium, vanadium, zirconium,manganese, scandium, lanthanum, and rare earth elements and alloysthereof. Preferred intermetallic precursors include aluminum, nickel,titanium, cobalt, silicon, iron, and refractory metals. Pluralintermetallic materials may, of course, be present in the matrix. It isnoted that the terms intermetallic containing matrix, or intermetallicmatrix, as used herein, are meant to define a matrix which ispredominantly intermetallic, although other materials, e.g., metalintermetallic precursors, may also be present in lesser amount.

Aluminides and silicides constitute the preferred intermetallic matrixmaterials of the present invention. Among the metallic elements capableof forming aluminides are titanium, nickel, iron, cobalt, and refractorymetals such as niobium, zirconium, tantalum and the like. Titanium formsthe aluminides Ti₃ Al, TiAl, and Al₃ Ti, while nickel forms Ni₃ Al andNiAl. Other aluminides include Zr₃ Al, Co₃ Al and Fe₃ Al. For thepurposes of the present invention, the aluminides of titanium and nickeland the silicides of titanium are presently preferred. Substitution ofone or more elements within an intermetallic compound is possible, andmay be desirable to effect crystal lattice changes. Exemplary is thesubstitution of aluminum by titanium in Ni₃ Al to form Ni₃ (Al, Ti). Inaddition, two phase mixtures such as TiAl and Ti₃ Al or TiAl and TiAl₃are suitable. Further, TiAl containing matrices may contain retainedbeta and ordered beta (B2).

It should be recognized that aluminides are not necessarily of acomposition such that the components are present as exact integers. Forexample, nickel aluminide is commonly referred to as Ni₃ Al although itis an intermetallic phase and not a simple ionic or covalently bondedcompound, as it exists over a range of compositions as a function oftemperature, e.g., about 72.5 to 77 weight percent nickel (85.1 to 87.8atomic percent) at about 600° C. Thus, aluminides, and intermetallicsgenerally, may be defined as the compounds which form uponsolidification of near stoichiometric amounts of the appropriate metals.In the molten state, however, the intermetallics exist primarily as arandom mixture of the elements thereof, possessing only relatively shortrange order. Within the scope of the present invention, this randommixture, or intermetallic derived liquid, may act as a solvent matrixmaterial through which the solvent assisted reaction of second phaseprecursors occurs. This molten state may thus be referred to as an"intermetallic derived solvent", or "solvent matrix material" whichterms also encompass the molten state of one or more precursors of saidintermetallic.

For the purpose of illustrating the various reaction modes that may beused to form a second phase dispersion within an intermetallic matrix,one is referred to the previously cited U.S. Pat. No. 4,774,052.

Varying amounts of ceramic may be incorporated into the compositematerial, depending upon the end use and the properties desired in theproduct. As previously noted, for dispersion strengthened materialshaving high modulus, one may utilize a preferred range of from about 10percent by volume to about 25 percent by volume. However, the ceramicvolume fraction may be varied considerably, so as to produce a compositewith the desired combination of properties, within the range of fromabout 0.5 percent by volume up to the point at which ductility issacrificed to an unacceptable extent. In contrast, cermet-likecomposites of up to about 95 percent or more by volume of ceramicmaterial in the aluminide containing matrix may be produced. Preferredranges for such materials will, of course, be dependent upon the desiredend use. It is possible to effectively tailor the composition to achievea spectrum of properties by controlling the proportions of the reactantand solvent materials.

As was previously stated, the first stage of the present inventionprovides for the formation of one or more finely dispersed precipitatesin a matrix of one or more intermetallic containing materials. It isimportant that the second phase precipitate material is not soluble in,or reactive with, the intermetallic derived solvent, while theconstituents of the second phase, individually, are at least sparinglysoluble in the intermetallic derived solvent. Thus, the exothermicdispersion reaction mechanism depends upon a certain amount of eachsecond phase forming constituent dissolving and diffusing in theintermetallic derived solvent, and while in solution (either liquid orsolid state), reacting exothermically to form the insoluble precipitaterapidly as a very fine particulate. The intermetallic derived solvent orsolvent matrix material provides a medium in which the reactive elementsmay diffuse and combine. Once the initial reaction has occurred, theheat released by the exothermic reaction causes additional diffusion ofreactive components in the solvent matrix material, and allows thereaction to proceed. During the initiation and reaction extremely hightemperatures may be achieved in very short periods of time. During thistime frame, essentially all of the reactive constituents in the solventmetal react to form the insoluble second phase, which immediatelyprecipitates.

The reaction mass may be subjected to conventional RS immediatelyfollowing the dispersoid forming reaction. Alternatively, the reactionmass may be solidified, alloyed, or further processed for subsequent RS.Techniques available for RS include melt spinning or atomization, toproduce in ribbon or droplet form a matrix metal having submicronceramic or intermetallic particles dispersed therein. Conventional RSprocessing is taught in J. R. Pickens et al, The Effect of RapidSolidification On The Microstructure And Properties Of Aluminum PowderMetallurgy Alloys. Rapid Solidification Processing, Principles andTechnologies III, pp. 150-170, Claitors (1982); and M. Cohen et al,Rapid Solidification Processing-An Outlook, Rapid SolidificationProcessing, Principles and Technologies, ii, pp. 1-23, Claitors (1980),hereby incorporated by reference. RS processing concerns liquid alloyssubjected to cooling rates on the order of about 104° C. per second to10⁸ ° C. per second. Several techniques are well established in thestate of the art to economically fabricate RS alloys as ribbons,filaments, wires, flakes or powders alone or in combination in largequantities.

The most common methods of consolidation of RS processing powders arehot isostatic pressing and hot extrusion. Superplasticity may beinvolved in some cases and, if so, it permits isothermal forging ofextruded billets into near-final shapes. Of course, the latter featureis also embodied in hot isostatic pressing.

Incremental solidification (otherwise called layer glazing) provides away of building up a three-dimensional shape by means of rapidlysolidified layers. The rapid solidification and consolidation can alsobe carried out concurrently. Whatever the method of consolidation, thestructure/property relationships stemming from the rapid solidificationwill depend upon the efficacy of the consolidation process as well as onthe final heat treatment.

The following examples illustrate the precipitation of fine articles ofa dispersoid to produce a composite having an aluminide containingmatrix.

EXAMPLE 1

An intermetallic-ceramic composite containing about 35 weight percenttitanium diboride particles dispersed in a matrix of titanium aluminide(Al₃ Ti) is prepared as follows. A well-blended mixture of 202.5 gramsof aluminum, 239.5 grams of titanium and 55.7 grams of boron is madefrom powders of the respective elements and the mixture thenisostatically compacted with a pressure of about 35,000 psi. The formedcompact is heated in an inconel retort and a reaction initiated at about660° C., causing melting of the compact. Upon removal from the retort,the compact is subjected to X-ray analysis which indicates the presenceof TiB₂ and Al₃ Ti with only trace amounts of the initial elements. AnSEM analysis indicates that the titanium diboride particles aresubmicron and dispersed in a titanium aluminide matrix. EDS analysis ofthe particles indicates that the particles are essentially pure titaniumdiboride.

EXAMPLE 2

An intermetallic-ceramic composite of titanium diboride particles in amatrix of titanium aluminide (Ti₃ Al) is prepared as follows. A mixtureof 67.5 grams of aluminum, 359.2 grams of titanium and 55.7 grams ofboron is thoroughly blended and the mixture then compacted and heated inthe manner of Example 1. The reaction temperature is observed to beabout 660° C. The resultant material upon solidification is a dispersionof fine particles of titanium diboride in a matrix of titanium aluminide(Ti₃ Al).

EXAMPLE 3

An intermetallic-ceramic composite containing 35 weight percent finetitanium diboride particles in a matrix of titanium aluminide (TiAl) isprepared as follows. A powdered mixture of about 117 grams powderedaluminum, about 328 grams titanium and about 56 grams of boron isprepared and mixed well to insure uniformity. The mixture is compactedand heated in the manner of Example 1 to yield a composite of finetitanium diboride particles in a matrix of TiAl. Analysis of thecomposite also reveals a minor amount of Ti₃ Al.

EXAMPLE 4

An intermetallic-ceramic composite of zirconium diboride particlesdispersed in titanium aluminide (Ti₃ Al) is prepared as follows. Amixture of 10.8 grams zirconium, 2.5 grams boron, 76.6 grams titaniumand 13.8 grams aluminum is thoroughly blended and then processed in amanner similar to that of Example 1 to yield the composite.

EXAMPLE 5

An intermetallic ceramic composite of zirconium diboride particlesdispersed in a matrix of nickel aluminide (Ni₃ Al) is prepared. Twomixtures, each containing 97 grams of zirconium, 23 grams of boron, 243grams of nickel and 37.2 grams of aluminum are blended thoroughly. Onemixture is heated to reaction initiation temperature in a resistanceheated furnace, and the other heated to reaction initiation temperatureby induction heating. The resultant composites each contain a smallamount of unreacted nickel. When subjected to fracturing forces, thesecomposites have a fracture surface which exhibits microvoid coalescence,which tends to indicate that the mode of fracture was a ductile one,consistent with the fine grain size of the aluminide matrix.

EXAMPLE 6

A mixture of nickel, aluminum, titanium, and boron in the stoichiometricproportions for the formation of nickel aluminide (Ni₃ Al) and titaniumdiboride (TiB₂), i.e., 10 percent by weight aluminum, 62 percent byweight nickel, 19 percent by weight titanium, and 9 percent by weightboron, is compacted to 40,000 pounds per square inch, and then heated ina furnace. Upon reaching 620° C., a rapid exotherm is noted, whichsubsequent analysis by X-ray diffraction and scanning electronmicroscopy identifies as resulting from the formation of titaniumdiboride particles in a nickel aluminide matrix. It is evident from thisexperiment that a ceramic phase, e.g., titanium diboride, could bedirectly precipitated in an intermetallic phase, e.g., nickel aluminide,provided the affinity of the ceramic-forming species for each other isgreater than either has for the two elements making up the intermetallicmatrix.

EXAMPLE 7

An intermetallic-ceramic composite of titanium diboride particlesdispersed in a matrix of nickel aluminide (Ni₃ Al) is prepared asfollows. A mixture of 103.5 grams titanium, 46.5 grams boron, 302.5grams nickel, and 47.5 grams aluminum is blended and then isostaticallypressed. About 100 grams of the pressed compact is then reacted in aretort. From a temperature probe placed adjacent to, but not touchingthe compact, the reaction apparently occurs at about 807° C. and thetemperature during reaction peaks at about 1050° C. X-ray diffraction ofthe resultant composite indicates the presence of TiB₂, Ni₃ Al, andresidual Ni.

EXAMPLE 8

In a series of experiments, the formation of each of the dispersoidshafnium carbide, zirconium carbide, titanium carbide, titanium boride,titanium diboride, and vanadium diboride in the matrices of titaniumaluminide (Ti₃ Al) and nickel aluminide (Ni₃ Al) is investigated. Inpreparing the various composites, the constituents forming the ceramicdispersoid and the components forming the aluminide containing matrixare reacted at the same time. The constituents and components in thereacting mixture are combined in such proportions so as to yield anintermetallic matrix composite containing about 40 weight percentceramic dispersoid. The reactions for each composite are conductedtwice, one of the reactions being conducted under an argon atmosphereand the other under vacuum. Induction heating is used to initiate eachreaction, and at the first indication of a reaction, power to theinduction heating unit is terminated so that the composite may cool asquickly as possible.

Upon completion of the reaction, each of the formed ceramic-aluminidecomposites is examined by X-ray diffraction analysis to determine itscomposition. In addition, a small amount of the matrix is dissolved inacid and the ceramic particles are observed for particle size by ascanning electron microscope and also examined by X-ray diffraction todetermine the particle composition.

The results of these observations are set forth in Table I.

                                      TABLE I                                     __________________________________________________________________________    Desired                     Reacted Under Vacuum                              Compound                                                                              Reacted under Argon                Particle                           Dispersoid/                                                                           Major   Minor  Particle                                                                           Major   Minor  Size                               Matrix  Component                                                                             Component                                                                            Size Component                                                                             Component                                                                            (micron)                                                                           Comment                       __________________________________________________________________________    HfC/Ni.sub.3 Al                                                                       HfC, Ni.sub.3 Al,                                                                     Hf, Al.sub.2 Hf                                                                      0.1-2                                                                              HfC, Ni.sub.3 Al                                                                      Hf     0.1-1.3                                                                            Amount of Al.sub.2 Hf                 Ni                  Ni                  was very small                HfC/Ti.sub.3 Al                                                                       HfC, TiC                                                                              Ti.sub.3 AlC                                                                         0.1-1.2                                                                            HfC, TiC                                                                              Ti.sub.3 AlC                                                                         0.5                                        Ti.sub.3 Al, Ti     Ti.sub.3 Al, AlTi.sub.2                           VB.sub.2 /Ni.sub.3 Al                                                                 VB.sub.2, Ni, Al,                                                                     --     0.1-1.3                                                                            VB.sub.2, Ni.sub.3 Al,                                                                V      1-2                                        V.sub.3 B.sub.4     V.sub.3 B.sub.4, Ni, Al                           VB.sub.2 /Ti.sub.3 Al                                                                 Ti.sub.3 Al,                                                                          V, TiB.sub.2                                                                         none Ti.sub.3 Al,                                                                          TiB.sub.2,                                                                           none                                       V.sub.3 V.sub.4     Al.sub.11 V                                                                           Al.sub.6 V                                        Al.sub.11 V,                                                                  Al.sub. 6 V                                                           TiC/Ni.sub.3 Al                                                                       TiC, Ni.sub.3 Al                                                                      --     0.1-2                                                                              TiC, Ni.sub.3 Al,                                                                     --     0.1-1.5                                    Ni                  Ni                                                TiC/Ti.sub.3 Al                                                                       TiC, Ti.sub.3 AlC                                                                     Ti.sub.3 Al                                                                          0.1-1.5                                                                            TiC, Ti.sub.3 AlC                                                                     --     0.2-2                              ZrC/Ni.sub.3 Al                                                                       ZrC, Ni.sub.3 Al,                                                                     Ni.sub.5 Zr                                                                          1    ZrC, Zr, Ni,                                                                          --     0.5  Intermetallics of                     Ni.sub.7 Zr.sub.2   Ni.sub.3 Al         Ni and Zr are                                                                 probably unstable                                                             phases                        ZrC/Ti.sub.3 Al                                                                       TiC, Ti.sub.3 Al,                                                                     AlZr.sub.3                                                                           none Ti.sub.3 AlC, TiC,                                                                    --     none Did not form ZrC                      Ti.sub.3 AlC        Al.sub.2 Zr         dispersoid, due                                                               to greater sta-                                                               bility of TiC,                                                                Ti.sub.3 AlC                  TiB/Ti.sub.3 Al                                                                       Ti.sub.3 Al                                                                           TiB    Ti.sub.2 Al                                                                        --      --     --                                 (10 vol %)                                                                    TiB/Ti.sub.3 Al                                                                       Ti.sub.3 Al                                                                           TiB    Ti.sub.2 Al                                                                        --      --     --                                 (30 vol %)                                                                    TiB and TiC/                                                                          TiC     --     --   --      --     --   Using B.sub.4 C as            Ti.sub.3 Al                                                                           TiB                                     reactant                              Ti.sub.3 Al                                                           TiB.sub.2 /TiAl                                                                       TiAl    TiB.sub.2                                                                            --   --      --     --                                 TiC/TiAl                                                                              TiC     Ti.sub.3 Al                                                                          --   --      --     --                                         TiAl    Ti.sub.2 Al                                                   __________________________________________________________________________

EXAMPLE 9

An intermetallic-ceramic composite having mixed ceramic dispersoids isprepared by mixing 11.0 grams of Al₄ C₃, 33.8 grams of tantalum, and135.2 grams of niobium, and heating in a graphite induction furnace.Analysis of the recovered product reveals the presence of both TaC andTa₂ C in a matrix of Nb₃ Al.

EXAMPLE 10

An intermetallic-ceramic composite comprising a ceramic dispersoid in amixed intermetallic matrix is prepared by mixing 43.6 grams of titanium,123.6 grams of tantalum, and 32.8 grams of A₁₄ C, compacting, andreacting on a water cooled copper holder in an induction furnace underflowing argon. Upon recovery of the reaction product, X-ray analysisshows the presence of TiC and a mixed matrix of TiAl, TaAl₃, and TaAl₂.

EXAMPLE 11

An intermetallic-ceramic composite of titanium diboride in a matrix oftitanium aluminide (Al₃ Ti) is prepared by the master concentrate route.A solidified melt comprising 30 weight percent titanium diboride in atitanium aluminide matrix is comminuted to particles having an averagesize of about 1 millimeter, and then a melt of about 860 grams oftitanium aluminide is prepared under a protective inert atmosphere. Theparticles are then added to the melt and held at that temperature for asufficient period of time to insure complete melting of theintermetallic phase of the particles, and a uniform distribution of thetitanium diboride. The melt is then solidified to yield a composite of15 weight percent titanium diboride dispersed in a titanium aluminidematrix.

EXAMPLE 12

An intermetallic-ceramic composite is prepared by the direct additionroute, by mixing 65.5 grams of titanium, 10 grams of boron, and 24.2grams of aluminum, compacting, and incrementally adding to a molten poolof TiAl under inert atmosphere. On addition of the compact to the moltenpool, a reaction occurs resulting in the formation of fine, evenlydispersed TiB₂ particles. Upon completion of the addition, the mixtureis cast and recovered as a dispersion of TiB₂ in a matrix of TiAl.

While it is believed that all of the composite materials produced by themethods exemplified above may be subjected to rapid solidification inaccordance with the present invention, the following examples set forthspecific experimental data relative to this invention.

Ingots of intermetallic-second phase composite material produced inaccordance with the present invention are converted into powder usingconventional inert gas centrifugal atomization techniques. Thesetechniques involve heating the ingot until melting of the intermetallicmatrix is complete, followed by rapid solidification of the melt to forma composite powder. Compositions of theintermetallic-second phasecomposites ingots atomized are given in Table II.

                  TABLE II                                                        ______________________________________                                        Matrix Composition                                                                            Dispersoid Loading                                                                          Dispersoid                                      (atomic percent)                                                                              (weight percent)                                                                            Composition                                     ______________________________________                                        Ti--45 Al       7.5           TiB.sub.2                                       Ti--45 Al       10.0          TiB.sub.2                                       Ti--45 Al       10.0          TiB.sub.2 /TiC                                  Ti--45 Al       7.5           TiB.sub.2 /TiN                                  Ti--48 Al       7.5           TiB.sub.2                                       Ti--48 Al--4 Nb 7.5           TiB.sub.2                                       Ti--48 Al--2 V  7.5           TiB.sub.2                                       Ti--45 Al--0.5 Er                                                                             7.5           TiB.sub.2                                       Ti--47 Al--6 Nb--0.5 Er                                                                       6.0           TiB.sub.2                                       Ti--47 Al--10 Nb--0.5 Er                                                                      6.0           TiB.sub.2                                       Ti--45 Al--0.5 Er                                                                             7.5           TiB.sub.2 /TiN                                  Ti--45 Al       4.5           NbB.sub.2                                       Ti--45 Al--5 Mo 9.0           NbB.sub.2                                       ______________________________________                                    

Upon chemical analysis of the as-atomized powder, it is noted thatmaterial containing TiB₂ in the precursor ingot was more resistant tocarbon contamination from the graphite melting crucible, relative toalloys which did not previously contain TiB₂. For example, a compositecomprising 7.5 volume percent TiB₂ particles dispersed within a Ti-45Almatrix produced in accordance with the methods of the present inventionis found to contain 0.116 weight percent carbon. A composite comprising7.5 volume percent TiB₂ particles dispersed within a Ti-48Al-4Nb matrixis found to contain 0.085 weight percent carbon, while a compositecomprising 7.5 volume percent TiB₂ particles dispersed within aTi-47Al-3Nb matrix is found to contain 0.125 weight percent carbon.

Upon microstructural examination of the as-atomized powder, severalfeatures unique to the RS products of the present invention areobserved.

The intermetallic matrices of rapidly solidified powders produced by themethods of the present invention have been found to be devoid ofdedritic structure which is otherwise present in conventional RSpowders. FIG. 2, which is a photomicrograph showing a cross section of aconventional unreinforced, rapidly solidified titanium aluminide(Ti-48Al) powder, illustrates the dendritic microstructure present inprior art RS powders. By comparision, no such dendritic microstructureis observed in the RS powders of the present invention (see FIG. 1).

In the case of RS powders comprising TiB₂ particulates within titaniumaluminide matrices, the presence of fine boride particles with aflake-like morphology is observed within the microstructure of eachpowder particle. Chemical dissolution of the matrix coupled withsubsequent x-ray diffraction identifies this flake-like RS phase asTiB₂. This phase apparently forms as a result of high undercooling andenhanced nucleation kinetics attributable to the rapid solidificationtechnique. It can thus be inferred that a higher solubility of boronexists in the molten state relative to that in the solid state. Inaddition, the presence of equiaxed TiB₂ particulates of comparable sizeand morphology to those present in the starting ingots is observed,indicating that a substantial portion of the pre-formed second phaseparticles have survived the remelting and subsequent RS atomizationprocess. Thus, in RS powders comprising TiB₂ particulates withintitanium aluminide matrices, a bimodal distribution is achievedcomprising equiaxed TiB₂ particles in the size range of about 1 to 5microns and finer flake-like TiB₂ particles. The widths of theseflake-like particles may range from about 0.05 to about 0.2 microns,while the lengths may range from about 0.5 to about 5 microns. FIG. 1illustrates the microstructure of an RS powder produced in accordancewith the present invention comprising a bimodal distribution of equiaxedand flake-like TiB₂ particulates within a Ti-48Al-4Nb matrix. FIG. 3illustrates a similar microstructure comprising a bimodal distributionof equiaxed and flake-like TiB₂ particulates within a Ti-45Al matrix.

Upon heat treatment of the titanium aluminide-TiB₂ RS powders producedby the present invention, the flake-like TiB₂ particles described aboveundergo a change in shape, with a reduction in aspect ratio approachingspherical. The resulting particulate distribution is superior to that ofthe precursor ingot material in terms of its reduced particle size,improved particle distribution, and increased number density. Theefficiency of the morphological transformation has been found toincrease with increasing heat treatment temperatures to approximately1125° C. Thus, after heat treatment, a bimodal distribution is achievedcomprising larger equiaxed TiB₂ particles in the size range of about 1to about 5 microns and finer equiaxed TiB₂ particles in the size rangeof about 0.1 to about 1 micron. FIG. 4 illustrates the microstructure ofa heat treated RS powder produced in accordance with the presentinvention comprising a bimodal distribution of equiaxed TiB₂particulates within a Ti-48Al-2V matrix.

It is understood that the above description of the present invention issusceptible to considerable modification, change, and adaptation bythose skilled in the art. Rapidly solidified products of the presentinvention may be used with prior art powder metallurgy techniques forforming composites, for example. The product of the present inventionmay also be subjected to subsequent heat treatment or combination withprior art products and structures, including filaments, and the like.Accordingly, all such modifications, changes, and adaptations areintended to be considered to be within the scope of the presentinvention, which is set forth by the appended claims.

We claim:
 1. A rapidly solidified composite powder comprising boride,carbide, nitride or silicide second phase particles dispersed within atitanium aluminide containing matrix.
 2. The rapidly solidifiedcomposite powder of claim 1, wherein the titanium aluminide is Ti₃ AI,TiAl or a combination thereof.
 3. The rapidly solidified compositepowder of claim 1, wherein the titanium aluminide is TiAl, TiAl₃ or acombination thereof.
 4. The rapidly solidified composite powder of claim1, wherein the second phase particles comprise from about 1 to about 40volume percent of the composite.
 5. The rapidly solidified compositepowder of claim 1, wherein the second phase comprises TiB₂.
 6. Therapidly solidified composite powder of claim 5, wherein the TiB₂particles comprise a bimodal distribution characterized by substantiallyequiaxed TiB₂ particles in a size range between about 1 and about 5microns and flake-like TiB₂ particles having widths between about 0.05and about 0.2 microns and lengths between about 0.5 and about 5 microns.7. The rapidly solidified composite powder of claim 6, wherein thetitanium aluminide containing matrix is substantially dendrite free. 8.The rapidly solidified composite powder of claim 5, wherein the titaniumaluminide is Ti₃ Al, TiAl or a combination thereof.
 9. The rapidlysolidified composite powder of claim 5, wherein the titanium aluminideis TiAl, TiAl₃ or a combination thereof.
 10. The rapidly solidifiedcomposite powder of claim 5, wherein the titanium aluminide is alloyedwith additional elements.
 11. The rapidly solidified composite powder ofclaim 5, wherein the TiB₂ particles comprise from about 0.5 to about 40volume percent of the composite.
 12. The rapidly solidified compositepowder of claim 5, wherein the TiB₂ particles are in-situ precipitated.13. The rapidly solidified composite powder of claim 1, furthercomprising a dispersion of rare earth oxide particulates within thetitanium aluminide containing matrix.
 14. The rapidly solidifiedcomposite powder of claim 13, wherein the rare earth oxide is erbiumoxide.
 15. A heat treated rapidly solidified composite materialcomprising TiB₂ second phase particles dispersed within a titaniumaluminide containing matrix wherein the TiBw particles comprise abimodal distribution characterized by substantially equiaxed TiB₂particles in a size range between about 1 and about 5 microns and finersubstantially equiaxed TiB₂ particles in a size range between about 0.1and about 1 micron.